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United States Patent |
5,273,709
|
Halverson
,   et al.
|
December 28, 1993
|
High neutron absorbing refractory compositions of matter and methods for
their manufacture
Abstract
Neutron absorbing refractory B.sub.4 C-Gd and Gd.sub.2 O.sub.3 -Gd cermets,
B.sub.4 C-Gd and Gd.sub.2 O.sub.3 -Gd metal-matrix composites, and B.sub.4
C-Gd.sub.2 O.sub.3 ceramic-ceramic composites can be manufactured by
applying fundamental thermodynamic and kinetic guidelines as processing
principals.
Three steps are involved in the fabrication of these new compositions of
matter. First, the starting materials are consolidated into a compacted
porous green body. Next, the green body is densified using the appropriate
method depending on the class of material sought: cermet, metal-matrix
composite, or ceramic-ceramic composite. Finally, either during the
densification process or by subsequent heat treatment, new phase evolution
is obtained via interfacial chemical reactions occurring in the
microstructures.
The existence of a new phase has been identified in B.sub.4 C-Gd and
B.sub.4 C-Gd.sub.2 O.sub.3 composites.
Inventors:
|
Halverson; Danny C. (Modesto, CA);
Billings; Garth W. (Auburn, CA);
Johnston; George M. (Santa Rosa, CA)
|
Assignee:
|
Thermal Technology Inc. (Santa Rosa, CA)
|
Appl. No.:
|
933859 |
Filed:
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August 24, 1992 |
Current U.S. Class: |
419/45; 419/14; 419/19; 419/20; 419/38 |
Intern'l Class: |
B22F 003/12 |
Field of Search: |
419/14,15,17,19,20,24,38,48,45
|
References Cited
U.S. Patent Documents
2992178 | Jul., 1961 | Lustman et al. | 252/478.
|
3245782 | Apr., 1966 | Ray | 252/478.
|
3356618 | Dec., 1967 | Rich et al. | 252/478.
|
4474728 | Oct., 1984 | Radford | 376/339.
|
4605440 | Aug., 1986 | Halverson et al. | 75/238.
|
4636480 | Jan., 1987 | Hillig | 501/87.
|
4671927 | Jun., 1987 | Alsop | 376/419.
|
4744922 | May., 1988 | Blakely et al. | 252/478.
|
4789520 | Dec., 1988 | Morimoto et al. | 376/419.
|
4826630 | May., 1989 | Radford et al. | 252/478.
|
Primary Examiner: Walsh; Donald P.
Assistant Examiner: Jenkins; Daniel
Attorney, Agent or Firm: Sartorio; Henry P.
Parent Case Text
This application is a division of application Ser. No. 07/591,094, filed
Oct. 1, 1990, now U.S. Pat. No. 5,156,804.
Claims
We claim:
1. A method of forming a neutron absorbing refractory composite material,
comprising:
selecting a pair of initial reactants from the group consisting of: (a)
gadolinium or alloys thereof, (b) boron carbide, and (c) gadolinium oxide;
consolidating the initial reactants;
densifying the consolidated initial reactants;
reacting the densified initial reactants to produce a material of
preselected composition having a plurality of interfacial reaction product
phases.
2. The method of claim 1 wherein the initial reactants are in the form of
particles, fibers, or whiskers.
3. The method of claim 1 wherein the step of consolidating the initial
reactants comprises:
selecting starting particle size distributions;
dispersing and mixing the particles;
consolidating the dispersed and mixed particles substantially
homogeneously.
4. The method of claim 1 for producing a B.sub.4 C-Gd or Gd.sub.2 O.sub.3
-Gd cermet, wherein the step of densifying the initial reactants comprises
wetting the B.sub.4 C or Gd.sub.2 O.sub.3 phase by the Gd metal phase.
5. The method of claim 1 for producing a B.sub.4 C-Gd or Gd.sub.2 O.sub.3
-Gd metal-matrix composite, wherein the step of densifying the initial
reactants comprises producing plastic flow of the Gd metal phase around
the B.sub.4 C or Gd.sub.2 O.sub.3.
6. The method of claim 1 for producing a B.sub.4 C-Gd.sub.2 O.sub.3
composite wherein the step of densifying comprises producing interphase
diffusion and rearrangement.
7. The method of claim 1 wherein the step of reacting is performed by
heating at a preselected temperature for a preselected time.
8. The method of claim 3 wherein the step of selecting starting particle
size distributions comprises selecting a plurality of different particle
sizes to enhance particle packing density.
9. The method of claim 3 wherein the step of dispersing and mixing is
performed by colloidal dispersion and ultrasonic mixing.
10. The method of claim 3 wherein the step of consolidating is selected
from:
(a) filtering followed by cold pressing,
(b) slip casting,
(c) pressure casting.
11. The method of claim 1 wherein the step of consolidating is performed by
injection molding or extruding the initial reactants.
12. The method of claim 1, for producing a B.sub.4 C-Gd cermet, further
comprising applying external pressure during the densification step.
13. The method of claim 12 comprising performing the step of densification
by squeeze casting, hot pressing, or hot isostatic pressing.
14. The method of claim 1, for producing a B.sub.4 C-Gd cermet, wherein the
densification step is performed at a temperature of about 1200.degree.
C.-1300.degree. C. for a time of less than about 10 minutes.
15. The method of claim 4, for producing a Gd.sub.2 O.sub.3 -Gd cermet,
wherein the densification step is performed by pressureless liquid-phase
sintering or liquid-metal infiltration.
16. The method of claim 5, for producing a B.sub.4 C-Gd metal-matrix
composite, wherein the densification step is performed at about the B-C-Gd
ternary eutectic temperature (1200.degree. C.).
17. The method of claim 5, for producing a Gd.sub.2 O.sub.3 -Gd
metal-matrix composite, wherein the densification step is performed at
about the liquidus temperature of Gd (1300.degree. C.).
18. The method of claim 5 wherein the densification step is performed by
vacuum hot pressing or induction melting.
19. The method of claim 6, for producing a Gd.sub.2 O.sub.3 -rich
composite, wherein the densification step is performed by solid-state
diffusion.
20. The method of claim 6, for producing a B.sub.4 C-rich composite,
wherein the densification step is performed by liquid-phase sintering.
21. The method of claim 20 further comprising applying external pressure
during the densification step.
22. The method of claim 21 wherein the densification step is performed by
hot pressing or hot isostatic pressing.
23. A method of forming a ceramic-metal composite material comprising:
selecting a ceramic phase from B.sub.4 C and Gd.sub.2 O.sub.3, and a metal
phase from Cd, In, Te, Pb, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb,
Lu, and alloys thereof;
consolidating the ceramic phase and metal phase;
densifying the consolidated ceramic and metal phases;
reacting the densified ceramic and metal phases to produce a plurality of
interfacial reaction product phases.
Description
BACKGROUND OF THE INVENTION
This invention relates generally to boron carbide-gadolinium oxide, boron
carbide-gadolinium, and gadolinium oxide-gadolinium compositions of matter
and more particularly to boron carbide-gadolinium oxide ceramic composites
and boron carbide-gadolinium and gadolinium oxide-gadolinium cermets or
metal-matrix composites.
U.S. Pat. No. 4,605,440 by Halverson, Pyzik and Aksay, U.S. Pat. No.
4,704,250 by Cline and Fulton, and U.S. Pat. No. 4,718,941 by Halverson
and Landingham all pertain to boron carbide-reactive metal composites and
their manufacture. However, these patents do not show specific boron
carbide-gadolinium compositions or methods for producing directly usable
consolidated bodies thereof. U.S. Pat. Nos. 4,826,630 and 4,474,728 by
Carlson and Radford, U.S. Pat. No. 4,744,922 by Blakely and Shaffer, U.S.
Pat. No. 4,671,927 by Alsop, and U.S. Pat. Nos. 4,657,876 and 4,636,480 by
Hillig describe various neutron absorbing materials.
The field of nuclear physics has matured and created a revolutionary impact
on modern history. The applications of research in neutron-induced
reactions appear both in other areas of fundamental research and in the
practical areas of nuclear energy production.
The ultimate fate of a free neutron liberated through some reaction is
either absorption by a nucleus, or transformation by beta decay. The
latter process is so weak as to be negligible for practical applications.
By a wide margin, the most important absorption process in a non-fissile
nucleus is radiative capture.
Neutron capture can occur over nine orders of energy magnitude, slow or
thermal neutrons from as low as 10.sup.-2 eV to fast neutrons as high as
14 MeV. Different atomic mechanisms are associated with the radiative
capture of each. Neutron capture processes have lifetimes ranging from
10.sup.-22 seconds (thermal capture) to as long as 10.sup.-15 seconds
(fast capture).
For practical applications involving slow neutrons, it is sufficient to
work with the statistical theory of radiative capture. This topic is
highly complex; however, qualitative consideration of materials can be
made by initially examining their average capture cross section
properties.
The nuclear cross section of a material is a measure of the probability of
a particular process. In the case of radioative capture, the capture cross
section is expressed as .sigma., and is the effective target area of the
nucleus with which a neutron must interact to produce a given reaction.
The unit of .sigma. is the barn (1 barn=10.sup.-24 cm.sup.2). Absorption
cross sections for thermal neutrons range from 4.6.times.10.sup.-4 barn
for deuterium to 3.3.times.10.sup.6 barns for xenon.sup.135.
To protect nuclear reactor operating personnel against damaging biological
effects of neutrons and gamma rays, shielding is required around nuclear
reactors. Neutron and gamma ray fluxes in the range of 10.sup.13 to
10.sup.14 must be attenuated to 10.sup.3 particles/cm.sup.2 /sec to meet
tolerance radiation levels.
To attenuate gamma rays, which interact primarily with the orbital
electrons of atoms, a material with high atomic number containing a high
density of these electrons is required. Examples are lead, tungsten,
depleted uranium, or concrete containing high-Z elements in the form of
scrap or heavy ore.
To attenuate neutrons, they must be slowed down and then absorbed.
Hydrogenous materials such as water, concrete, or polyethylene are
excellent moderators. The slowed neutrons must be absorbed without
producing high-energy-capture gamma rays. This has historically been
accomplished by using boron.sup.10 ; however, even with boron.sup.10 some
gamma shielding outside the neutron shield is generally required.
When a neutron is absorbed by the nucleus of an atom an exothermic process
results and the compound nucleus reaches an excited energy state between 4
and 10 MeV, as determined by the center-of-mass kinetic energy and the
rest-mass energy difference of the final and initial nuclides. This state
decays by the emission of electromagnetic (gamma) radiation, leaving the
compound nucleus in a lower energy state. Subsequent radiative decay of
this and lower energy states, i.e., a cascade of gamma rays, leaves the
compound nucleus in its ground state, which may, or may not, be stable
against alpha or beta decay (daughter products).
The inherent atomic processes associated with the radiative capture of
neutrons results in exothermic reactions that typically prohibit the use
of hydrogenous materials because of their low-temperature phase changes;
e.g., water boils at 100.degree. C., polyethylene softens near 87.degree.
C.
Boron metal also has its drawbacks. It corrodes easily and is physically
unstable under irradiation. Alloying to overcome these problems merely
reduces the boron content of the absorbing material. Because of these
concerns, boron carbide has been used extensively as a neutron absorbing
material in various types of nuclear reactors for several decades.
Boron carbide exists as a homogeneous range of boron and carbon
compositions between 9 and 24 at. % C. The most common stoichiometries
being B.sub.4 C (B.sub.12 C.sub.3) and B.sub.13 C.sub.2, both of which are
boron rich. Richer boron stoichiometries, B.sub.8 C and B.sub.25 C, are
also known to exist; however, these are less favored thermodynamically.
The high boron content and refractory nature of B.sub.4 C (melting point
.apprxeq.2350.degree. C.) made it a choice candidate for high temperature
neutron absorbing reactions.
The ideal neutron absorbing material would be light weight, refractory, not
impart long-lived daughter products, be thermally shock resistant, of low
density yet not too porous, be resistant to corrosion and oxidation, have
high fracture toughness and high strength, and not promulgate dust on
delivery or while in use. Low cost, obviously, would be another attractive
advantage.
Boron carbide is refractory, has a specific gravity of 2.52, a modulus of
rupture .apprxeq.300 MPa, and can be hot pressed into fully dense bodies.
Boron carbide displays low fracture toughness, and also rapidly oxidizes
above 800.degree. C. In addition, boron carbide's thermal shock resistance
is poor.
One way to increase the fracture toughness and thermal shock resistance of
B.sub.4 C is by the addition of a metal phase, e.g., B.sub.4 C-metal
cermets or metal-matrix composites. A cermet is defined as a ceramic-metal
composite such that the final microstructure is .gtoreq.50 vol. % ceramic
phases. A metal matrix composite is defined as a ceramic-metal composite
such that the final microstructure is <50 vol. % ceramic phases. The
ceramic phases can be the initial starting ceramic materials or reaction
products that result from chemical reactions between two ceramic phases or
between ceramic and metal phases.
Another way to increase the fracture toughness and strength of B.sub.4 C is
through the introduction of another ceramic phase, e.g., B.sub.4
C-Al.sub.2 O.sub.3, B.sub.4 C-TiB.sub.2, and B.sub.4 C-SiC composites.
Although large increases in fracture toughness and strength are generally
obtained with the addition of a metal phase, the introduction of another
ceramic phase can increase toughness while maintaining the refractory
nature of the composite.
One of the most appropriate metal phases to consider for the absorption of
neutrons is gadolinium, Gd. This metal also exists in the form of a stable
oxide known as gadolinium oxide, Gd.sub.2 O.sub.3. Gadolinium has the
highest nuclear capture cross section of any element known,
.sigma..apprxeq.40,000 barns, compared to B.sup.10 with a
.sigma..apprxeq.4,000 barns.
Gadolinium offers mechanical and physical properties conducive to
fabricating B.sub.4 C-Gd cermets or metal-matrix composites, according to
the invention, that approach "ideal" neutron absorbing material
conditions. For example, Gd is used as a burnable poison in shields and
control rods in nuclear reactors. It has a melting point of 1313.degree.
C., a boiling point of .apprxeq.3000.degree. C., and an
.alpha..fwdarw..beta. transformation temperature of 1235.degree. C.
Gadolinium tarnishes slightly in air at room temperature; however, even at
1000.degree. C. the oxidation rate is slow because of the formation of the
tightly adhering oxide on the surface. It does not react with cold or hot
water, but will react vigorously with dilute acids.
Gadolinium has the following mechanical properties: Tensile strength
.apprxeq.122 MPa, yield strength .apprxeq.17 MPa, elongation .apprxeq.47%,
reduction in area .apprxeq.58%, and an elastic modulus .apprxeq.56 GPa. It
also has a thermal expansion coefficient of .apprxeq.9.times.10.sup.-6
/.degree. C.
Gadolinium's very low modulus of elasticity indicates it should be
substantially more resistant to thermal shock than B.sub.4 C. By forming a
B.sub.4 C-Gd composite, according to the invention, it should be possible
to obtain a refractory body with very high neutron absorbing capability
and good thermal shock resistance. Because the specific gravity of Gd is
7.90, the addition of B.sub.4 C will also reduce the composite's weight
substantially. Also, according to the invention, the reactions between Gd
and B.sub.4 C during processing will introduce other ceramic phases into
the composite resulting in a higher overall fracture toughness.
According to the invention, similar material properties should be obtained
by combining Gd.sub.2 O.sub.3 and B.sub.4 C, or Gd.sub.2 O.sub.3 and Gd,
to form ceramic-ceramic composites, cermets, or metal-matrix composites.
For example, Gd.sub.2 O.sub.3 has a specific gravity of 7.41 and a elastic
modulus of .apprxeq.130 GPa.
Accordingly, it is an object of the present invention to provide
boron-carbide-gadolinium cermet compositions, boron-carbide-gadolinium
metal-matrix compositions, boron-carbide-gadolinium-oxide compositions,
gadolinium-oxide-gadolinium cermet compositions, and
gadolinium-oxide-gadolinium metal-matrix compositions.
It is also an object of the invention to provide methods for forming
boron-carbide-gadolinium cermet compositions, boron-carbide-gadolinium
metal-matrix compositions, boron-carbide-gadolinium-oxide compositions,
gadolinium-oxide-gadolinium cermet compositions, and
gadolinium-oxide-gadolinium metal-matrix compositions.
It is another object of the invention to provide boron-carbide-gadolinium
cermet compositions, boron-carbide-gadolinium metal-matrix compositions,
boron-carbide-gadolinium-oxide compositions, gadolinium-oxide-gadolinium
cermet compositions, and gadolinium-oxide-gadolinium metal-matrix
compositions with refractory microstructures, and methods for forming
same.
It is a further object of the invention to provide boron-carbide-gadolinium
cermet compositions, boron-carbide-gadolinium metal-matrix compositions,
boron-carbide-gadolinium-oxide compositions, gadolinium-oxide-gadolinium
cermet compositions, and gadolinium-oxide-gadolinium metal-matrix
compositions which are fully dense, and methods for forming same.
It is another object of the invention to provide articles of manufacture
made from boron-carbide-gadolinium cermet compositions,
boron-carbide-gadolinium metal-matrix compositions,
boron-carbide-gadolinium-oxide compositions, gadolinium-oxide-gadolinium
cermet compositions, and gadolinium-oxide-gadolinium metal-matrix
compositions.
It is also an object of the invention to provide methods for making
boron-carbide-gadolinium cermet compositions, boron-carbide-gadolinium
metal-matrix compositions, boron-carbide-gadolinium-oxide compositions,
gadolinium-oxide-gadolinium cermet compositions, and
gadolinium-oxide-gadolinium metal-matrix compositions, and articles of
manufacture thereof at relatively low cost.
SUMMARY OF THE INVENTION
The present invention provides a whole spectrum of specific compositions of
boron-carbide-gadolinium, boron-carbide-gadolinium-oxide, and
gadolinium-oxide-gadolinium composites which apply basic thermodynamic and
kinetic principles to achieve these compositions. The invention includes a
plurality of multiphase compositions, including fully dense
microstructures, and methods for selectively producing the desired
compositions.
According to the invention, there are three major steps in the formation of
these compositions of matter. First, the initial reactants must be
properly prepared. This is accomplished by selecting the appropriate
starting particle size distributions, dispersing and mixing the particles,
and consolidating the particles into a state that is ready for step two.
Second, the capillarity thermodynamic criteria of achieving a rapid
consolidation through the wetting of the B.sub.4 C or Gd.sub.2 O.sub.3
phase by the Gd metal phase (Gd metal or Gd alloy) must be obtained in the
case of the B.sub.4 C-Gd cermets and Gd.sub.2 O.sub.3 -Gd cermets,
respectively. Or, the condition of plastic flow of the Gd metal (or alloy)
around the B.sub.4 C or Gd.sub.2 O.sub.3 must be obtained in the case of
B.sub.4 C-Gd metal-matrix composites and Gd.sub.2 O.sub.3 -Gd metal-matrix
composites, respectively. Or, the solid-state and/or liquid-state
rearrangement of oxide and carbide phases around each other must be
obtained in the case of B.sub.4 C-Gd.sub.2 O.sub.3 composites.
The third step is to apply reaction thermodynamic criteria to the
boron-carbide-gadolinium, boron-carbide-gadolinium-oxide or
gadolinium-oxide-gadolinium compositions in order to achieve desired
reaction products in the microstructure. Through this step, it is possible
to take each respective composition of step two and react them to specific
end products which result in different microstructures than those obtained
in step two. In the case of the cermet or metal-matrix compositions, it is
also possible to completely react all of the Gd metal or alloy thereof and
any metastable phases which form during these processes to achieve a
composite material which is completely without any metal phase or any
phase representative of the initial starting constituents. That is, it is
possible to start with a cermet or metal-matrix composition and end up
with a multiphase ceramic composite.
In general it is necessary to apply the kinetics of how these
boron-carbide-gadolinium, boron-carbide-gadolinium-oxide, and
gadolinium-oxide-gadolinium compositions consolidate during the above
processes in order to select the appropriate method of manufacture. Final
consolidation involves the application of temperature to these bodies such
that the microstructural phases will flow together or sinter. It may also
involve the application of pressure with temperature in order to assure
that fully dense final products are obtained.
BRIEF DESCRIPTION OF THE DRAWINGS
FIG. 1 is a graph of neutron capture cross section versus gadolinium
content in boron-carbide-gadolinium composites.
FIGS. 2A and 2B illustrate the wetting step in accordance with the
invention.
FIG. 3 is a graph of contact angle data for gadolinium metal in a
boron-carbide substrate.
FIG. 4 is a vapor pressure curve for gadolinium.
FIG. 5 is a flow chart of the fabrication sequence for producing
boron-carbide-gadolinium and gadolinium-oxide-gadolinium metal-matrix
composites.
FIG. 6 is the phase equilibrium diagram for the GdBO.sub.3 -B.sub.2 O.sub.3
pseudooinary system.
FIG. 7 is a microstructure of a boron-carbide-gadolinium cermet.
FIG. 8 is a microstructure of a boron-carbide-gadolinium-oxide composite.
FIG. 9 is a microstructure of a gadolinium-oxide-gadolinium cermet.
DETAILED DESCRIPTION OF THE INVENTION
Gadolinium (or an alloy thereof) is a compatible metal phase with B.sub.4 C
or Gd.sub.2 O.sub.3 ceramics and Gd.sub.2 O.sub.3 is a compatible ceramic
phase with B.sub.4 C ceramic for the development of B.sub.4 C-Gd, B.sub.4
C-Gd.sub.2 O.sub.3, and Gd.sub.2 O.sub.3 -Gd refractory composites for
neutron absorption because the starting phases are reactive with each
other.
Gadolinium and its oxide are terrestrially stable metal and ceramic phases,
respectively. Gadolinium possesses the highest neutron capture cross
section of any element known. This metal and its oxide is the third most
abundant of the lanthanide series of elements in the periodic table. Only
cerium and samarium are more plentiful. Its abundance makes it a cost
effective alternative to boron-based neutron absorbing materials.
Combining Gd or Gd.sub.2 O.sub.3 with B.sub.4 C, or combining Gd.sub.2
O.sub.3 with Gd allows the useful nuclear, thermal, and mechanical
properties of each phase to act together to form refractory composites for
practical nuclear applications. In particular, these composites offer
properties that can be controlled through various processing routes to
obtain high neutron capture ability, reduced weight, thermal shock
resistance, improved fracture toughness, corrosion and oxidation
resistance. Calculated neutron absorption values (minimum and maximum) for
B.sub.4 C-Gd cermets and B.sub.4 C-Gd metal-matrix composites are given in
FIG. 1.
Potential applications of B.sub.4 C-Gd, B.sub.4 C-Gd.sub.2 O.sub.3, and
Gd.sub.2 O.sub.3 -Gd refractory composites include, but are not limited
to, neutron shielding, control rods, burnable absorbers, secondary
shutdown systems, spent fuel storage containers, transportation casks,
personnel protection, protection of electronics, and other waste handling
or encapsulation applications requiring neutron attenuation. These
composites can also serve as cathodes and anodes in long-life electronic
devices and laser systems.
The methods for forming B.sub.4 C-Gd, B.sub.4 C-Gd.sub.2 O.sub.3, and
Gd.sub.2 O.sub.3 -Gd refractory composites are described in the following.
The methods include three principal steps, including (1) consolidation or
preparation of the starting materials, (2) densification by producing the
right capillarity-thermodynamic condition in the case of systems processed
above the melting point of Gd, or by producing the correct plastic flow
conditions in the case of systems processed below the melting point of Gd,
or by producing the correct interphase diffusion and rearrangement
conditions in the case of processing with non-metallic phases, (3)
reacting the starting materials to produce the desired compositions.
CONSOLIDATION OF STARTING MATERIALS
Correct preparation of the starting materials is required to produce fully
dense microstructures or at least microstructures with negligible
porosity. Preparation and consolidation of the starting materials involves
three steps: Selection of the appropriate starting particle size
distributions, dispersion and uniform mixing of the appropriate particles,
and consolidation of these particles in a very homogeneous manner.
Selecting the correct particle size distributions of the starting materials
is important because it directly affects densification and reaction
kinetics. That is, large particles will have much less surface area than
small particles, making the available surface for chemical reactions
smaller. In addition, the selection of several different particle size
distributions in combination can enhance the particle packing density
making microstructural rearrangement distances substantially smaller
thereby promoting densification and reaction product formation.
Uniform mixing of the starting constituents can be performed by mechanical
mixers or vibratory shakers; however, optimum uniformity is usually
obtained through the use of colloidal techniques. Colloidal mixing
involves the dispersion of the starting particles in a compatible liquid
medium. The dispersion may be electrostatic, steric, or a combination of
the two depending on the surface characteristics of the particles being
dispersed. Mixing is then accomplished by ultrasonication of the
particle-fluid slip. Other mixing techniques may be used but they are
generally not as effective as ultrasonic methods.
The final step in controlling the packing morphology of the green or
prefired body involves the actual consolidation of the particles into a
desired shape. To do this, a method for removing the dispersion fluid must
be used. This can be as simple as filtering out the solid particles prior
to cold pressing them to shape, or it may involve the use of slip casting
or pressure casting techniques where the fluid is sucked or pushed out of
the slip, respectively. Other, more advanced methods, such as injection
molding or extrusion of the starting constituents may also be used.
DENSIFICATION
Once the conditions for achieving the optimum packing morphology have been
obtained, the second step, in the case of systems processed above the
melting point of Gd (e.g., B.sub.4 C-Gd cermets and Gd.sub.2 O.sub.3 -Gd
cermets), is to apply the capillarity thermodynamic criteria of achieving
rapid densification through the kinetics of microstructural rearrangement
via liquid-phase sintering. Due to a greater degree of reactivity, it is
more difficult to obtain this rapid consolidation in the case of B.sub.4
C-Gd cermets than it is in the case of Gd.sub.2 O.sub.3 -Gd cermets, and
the former requires the application of pressure with temperature to
externally accelerate the densification kinetics.
In Gd.sub.2 O.sub.3 -Gd cermets, the criterion of a low contact angle of
the liquid Gd on the solid Gd.sub.2 O.sub.3 must be achieved. This
condition is often referred to as wetting. In B.sub.4 C-Gd cermets, the
external application of pressure negates this requirement; however, any
wetting of solid B.sub.4 C by liquid Gd will assist the rearrangement
process during pressing.
Wetting is defined as any process in which a solid-liquid interface is
formed such that the measured angle through the molten liquid phase is
acute. This is illustrated in FIG. 2A. Non-wetting is illustrated in FIG.
2B. The driving force for wetting is a reduction in free energy of the
system, where the system is defined as the local solid, liquid, and vapor
phases that coexist.
During wetting in B.sub.4 C-Gd and Gd.sub.2 O.sub.3 -Gd cermets, chemical
nonequilibrium thermodynamic conditions exist between the solid, liquid,
and vapor phases of the system. This nonequilibrium manifests itself in
the form of interfacial reactions that continue until a state of chemical
equilibrium is reached in the system.
Interfacial chemical reactions result in mass transfer across the B.sub.4
C-Gd or Gd.sub.2 O.sub.3 -Gd interfaces. Mass transfer results in a net
decrease in the system free energy and usually begins and continues during
sintering until chemical equilibrium is achieved. This process ultimately
results in the formation of interfacial reaction products.
Contact-angle measurements can be made to quantify this wetting phenomenon.
This is easily done by heating the Gd metal atop a polished substrate of
B.sub.4 C or Gd.sub.2 O.sub.3. The contact angle is then measured in-situ
and recorded.
As an example, contact angle data for Gd on a B.sub.4 C substrate is
presented in FIG. 3 which shows the contact angle as a function of
temperature in a vacuum environment of .apprxeq.10.sup.-4 torr. It is also
important to consider the time variable in making these measurements.
Without the three coupled variables; temperature, time, and pressure, it
is not possible to accurately replicate the conditions necessary to
achieve a desired wetting condition. Processing temperatures for producing
B.sub.4 C-Gd cermets, under wetting conditions alone, are in the range of
just above 1200.degree. C. to 1300.degree. C. Processing times, due to the
very reactive nature of this system, are typically less than 10 minutes.
It is interesting to note that wetting in the B.sub.4 C-Gd system occurs
below the 1313.degree. C. melting point of Gd. This must be due to
existence of a B-C-Gd ternary eutectic. There is no published phase
equilibria for this system; however, a 1180.degree. C. binary eutectic is
reported for the B-Gd system indicating a high-probability for this
conclusion.
The time parameter is particularly important and the one that constrains
the liquid-phase sintering of B.sub.4 C-Gd and Gd.sub.2 O.sub.3 -Gd
cermets the most. This is because mass transfer across the interfaces in
these cermets is time dependent. The higher degree of reactivity in a
B.sub.4 C-Gd cermet indicates that mass transfer across its interface is
greater/faster than mass transfer across Gd.sub.2 O.sub.3 -Gd interfaces.
The application of external pressure to the sintering process using hot
pressing or hot isostatic pressing changes the time scale for
rearrangement and can allow densification to occur much quicker.
The parameter of atmosphere is also important because it affects the vapor
phase in the solid-liquid-vapor system. This becomes a significant issue
when processing molten Gd at pressures below its vapor pressure. Published
vapor pressure data for Gd is plotted in FIG. 4 and indicates that vacuum
processing of material systems containing Gd could occur without
evaporation of molten Gd at temperatures between 1200.degree.-1300.degree.
C. (1473-1573K) when the atmospheric pressure remains at or above
.apprxeq.10.sup.-2 torr. Actual vacuum processing of systems containing
molten Gd have shown that the kinetics of evaporation for Gd are slow and
that processing at 10.sup.-4 torr is possible with minimal contamination
to the furnace.
The key point of the second step of this invention, as applied to B.sub.4
C-Gd and Gd.sub.2 O.sub.3 -Gd cermets, is that highly reactive systems
like B.sub.4 C-Gd require the application of external pressure to achieve
densification. This implies that squeeze casting, hot pressing or hot
isostatic pressing will be required to densify these compositions. In less
reactive systems like Gd.sub.2 O.sub.3 -Gd, the application of pressure is
not required and complete densification can be achieved by pressureless
sintering or liquid-metal infiltration techniques.
In the case of systems processed below the melting point of Gd (e.g.,
B.sub.4 C-Gd or Gd.sub.2 O.sub.3 -Gd metal-matrix composites), the second
step of the invention is densification by establishing the correct plastic
flow conditions such that the Gd metal will flow around the B.sub.4 C or
Gd.sub.2 O.sub.3 ceramic phases.
The whole process is dependent on the work hardening, recovery,
recrystallization, and grain growth of Gd metal in combination with
particles, whiskers, or fibers of B.sub.4 C or Gd.sub.2 O.sub.3. A flow
chart of the fabrication sequence for metal-matrix composites is given in
FIG. 5.
B.sub.4 C or Gd.sub.2 O.sub.3 material is mixed with atomized Gd powder in
the desired proportions either dry, or wet (colloidal) and then dried. The
mixture is then compacted at room temperature using axial or isostatic
cold pressing methods. This compact is either vacuum hot pressed or
induction melted into a billet and the billet subsequently extruded.
Vacuum hot pressing or induction melting is an operation where the
metal+ceramic billet is taken to near the B-C-Gd ternary eutectic
temperature (.apprxeq.1200.degree. C.) in the case of B.sub.4 C-Gd
metal-matrix composites, or to near the liquidus of Gd
(.apprxeq.1300.degree. C.) in the case of Gd.sub.2 O.sub.3 -Gd
metal-matrix composites.
In the case of vacuum hot pressing, once the billet is pressed to full
density, the die is then cooled to a temperature below the solidus of Gd
(approximately 100.degree. C. below the liquidus or eutectic temperature)
while maintaining pressure and vacuum, the billet removed from the die,
and then from the hot press. In the case of induction melting, the
metal+ceramic mixture is heated in a vacuum induction furnace to at least
the melting point of Gd, inductively mixed in the molten state, allowed to
cool, and then removed from the furnace and crucible.
After vacuum hot pressing or induction melting the billet is extruded
through a die at a temperature .apprxeq.75% of the eutectic or liquidus
temperature. Other conventional metal working procedures may also be
carried out on the billets. The extrusion of tubes, rods, plates, bars,
and shaped sections are possible. Forging, rolling, and squeeze casting
operations are also feasible with B.sub.4 C-Gd and Gd.sub.2 O.sub.3 -Gd
metal-matrix composites.
Extrusion, forging, rolling, and casting pressures vary widely depending on
the process, but more importantly depending on the ceramic solids loading
of the composite being worked with, higher pressures being required the
greater the ceramic content. Nominal pressures are used with ceramic
loadings up to 25 vol. %. Substantially higher pressures are required up
through 40 vol. % ceramic constituent and subsequent hot isostatic
pressing may be required to densify metal matrix composites with initial
ceramic contents between 40-50 vol. %.
The key point of the second step of this invention, as applied to B.sub.4
C-Gd and Gd.sub.2 O.sub.3 -Gd metal-matrix composites, is that highly
reactive systems like B.sub.4 C-Gd require consolidation at temperatures
near the B-C-Gd ternary eutectic temperature in order to maximize the
plastic flow of the Gd matrix material around the B.sub.4 C particles,
fibers, or whiskers. In less reactive systems like Gd.sub.2 O.sub.3 -Gd,
the consolidation temperature must be closer to the liquidus of the Gd
matrix for optimizing the plastic flow around Gd.sub.2 O.sub.3 particles,
fibers, or whiskers.
In the case of systems processed from refractory starting constituents
(e.g., B.sub.4 C-Gd.sub.2 O.sub.3 composites), the second step of this
invention involves producing the correct interphase diffusion and
rearrangement conditions such that the composite can achieve a fully dense
state.
This is accomplished primarily by solid-state diffusion; however, viscous
flow and liquid diffusion are also potential contributors to the process.
Solid-state sintering occurs readily between Gd.sub.2 O.sub.3 grains due
primarily to the ionic nature of this phase. B.sub.4 C, on the other hand,
does not easily sinter due to its covalent nature and low volume and grain
boundary diffusion rates. Hence, B.sub.4 C-Gd.sub.2 O.sub.3 systems that
are Gd.sub.2 O.sub.3 rich should be easier to sinter than systems that are
B.sub.4 C rich. This latter category, as well as the former, can be
densified using liquid-phase sintering techniques.
Sintering in the presence of a liquid phase is a way of bonding two or more
materials, which have different melting points, into dense bodies without
complete fusion of the more refractory Gd.sub.2 O.sub.3 or B.sub.4 C
phases. This type of sintering is unique in that the system remains
multiphase throughout the entire process and the maximum temperature
attained is between the liquidus and solidus of the system. The previously
discussed liquid-phase sintering of cermets showed how Gd was used as the
liquid phase. In the case of B.sub.4 C-Gd.sub.2 O.sub.3 compositions, the
Gd.sub.2 O.sub.3 phase or another phase becomes the liquid.
There are basically two types of sintering that can occur in the presence
of a liquid phase. The first type, liquid-phase sintering, uses a material
transport mechanism involving viscous flow and diffusion. The other type,
reactive liquid sintering, uses a material transport mechanism of viscous
flow and solution precipitation. Both of these mechanisms have the same
driving force which evolves from capillary pressures and surface tensions
occurring within the composite during sintering.
It is widely accepted that the sintering of nonoxide ceramics, like B.sub.4
C, to theoretical density is only possible with the aid of certain
additions of impurities which result in the formation of a liquid phase.
Without such additives it is necessary to apply pressure to the system in
order to force it to densify.
In Gd.sub.2 O.sub.3 -rich B.sub.4 C-Gd.sub.2 O.sub.3 composites the primary
driving force for densification is the difference in free energy or
chemical potential between the free surfaces of Gd.sub.2 O.sub.3 particles
and the points of contact between adjacent Gd.sub.2 O.sub.3 particles.
This solid-state diffusion process between the Gd.sub.2 O.sub.3 particles
involves material transport by surface and volume diffusion. Surface
diffusion does not result in densification; however, volume diffusion
does. Volume diffusion occurs along grain boundaries and through lattice
dislocations. Consequently, a system that is rich in Gd.sub.2 O.sub.3
should sinter to near full density without the application of external
pressure to the system. In addition, the reaction products that form
between Gd.sub.2 O.sub.3 and B.sub.4 C (e.g., GdBO.sub.3, GdB.sub.3
O.sub.6, B.sub.2 O.sub.3, etc.) may also serve as sintering aids to
densification.
In B.sub.4 C-rich B.sub.4 C-Gd.sub.2 O.sub.3 composites, just the opposite
occurs. There are not enough free surfaces and points of contact with
respect to the Gd.sub.2 O.sub.3 phase and consequently densification is
severely inhibited. This is quickly overcome, however, by use of hot
pressing or hot isostatic pressing methods.
Gd.sub.2 O.sub.3 melts at .apprxeq.2320.degree. C., forms a gadolinium
borate (GdBO.sub.3) that melts at .apprxeq.1590 .C, and forms a binary
eutectic with B.sub.2 O.sub.3 at .apprxeq.1230.degree. C. B.sub.4 C, on
the other hand, melts at .apprxeq.2350.degree. C. but can form a B203
surface layer in the presence of oxygen at much lower temperatures. This
shared characteristic between Gd.sub.2 O.sub.3 and B.sub.4 C assists in
the densification of B.sub.4 C-Gd.sub.2 O.sub.3 composites by providing a
low temperature sintering aid for the system. FIG. 6 shows the phase
equilibria for the GdBO.sub.3 -B.sub.2 O.sub.3 pseudobinary system. The
formation of a low melting point borate and eutectic clearly indicates
that liquid-phase densification is possible in B.sub.4 C-Gd.sub.2 O.sub.3
composites.
The key point of the second step of this invention, as applied to B.sub.4
C-Gd.sub.2 O.sub.3 composites, is that Gd.sub.2 O.sub.3 -rich systems do
not require the application of external pressure for densification while
B.sub.4 C-rich systems do.
Reaction Thermodynamics
If molten-metal liquid-phase sintering is to occur, the B.sub.4 C-Gd or
Gd.sub.2 O.sub.3 -Gd mixture must satisfy the reaction-thermodynamic
criterion that the solid B.sub.4 C phase and any metastable
gadolinium-borocarbide, gadolinium-boride, gadolinium-carbide compounds or
solid solutions be partially soluble in the liquid Gd or alloy phases
present; or that the solid Gd.sub.2 O.sub.3 phase and any metastable
gadolinium-oxide compounds or solid solutions be partially soluble in the
liquid Gd or alloy phases present.
If solid-state sintering and/or liquid-phase sintering is to occur in
B.sub.4 C-Gd.sub.2 O.sub.3 composites, then the B.sub.4 C-Gd.sub.2 O.sub.3
mixture must satisfy the reaction-thermodynamic criterion that the solid
B.sub.4 C and Gd.sub.2 O.sub.3 phases and any metastable
gadolinium-borocarbonate, gadolinium-borate, gadolinium-carbonate,
gadolinium-borocarbide, gadolinium-boride, adolinium-carbide,
gadolinium-oxide compounds or solid solutions be partially soluble in the
liquid GdBO.sub.3, liquid B.sub.2 O.sub.3, or other liquid phases present.
To fully understand how a particular B.sub.4 C-Gd, Gd.sub.2 O.sub.3 -Gd, or
B.sub.4 C-Gd.sub.2 O.sub.3 composite will react at different processing
isotherms, one must consider both thermodynamic and kinetic issues for
each respective composition. Merely examining phase-equilibria data is
often times not enough. This is because reactive compositions typically
form metastable interfacial phases. In many cases, the phase diagrams are
not available anyway.
Hence, detailed studies using x-ray diffraction and optical metallographic
equipment must be employed. The studies on each respective system need to
be done in incremental steps, evaluating the composition as a function of
processing history. Only in this way, can metastable phases and
equilibrium reaction products be correctly identified in a detailed
fashion.
In a more general sense, however, its is possible to determine the range of
compositions that are obtained from each respective system. The
composition ranges obtained for the B.sub.4 C-Gd, Gd.sub.2 O.sub.3 -Gd,
and B.sub.4 C-Gd.sub.2 O.sub.3 composites of this invention are presented
in Tables I-III.
TABLE I. Semi-quantitative x-ray diffraction analysis of B.sub.4 C-Gd
composites.
For Cermets (measured)
Major phase: New unknown phase(s)
Secondary phases: GdB.sub.4, GdC.sub.2, B.sub.8 C(?)
Minor phases: GdBC, Gd.sub.2 C.sub.3, Gd, B.sub.4 C, B.sub.13 C.sub.2,
GdB.sub.66, free boron(?)
Trace phases: Gd.sub.2 B.sub.5, B.sub.25 C, GdB.sub.2 (?), GdB.sub.6 (?),
C(?)
For Metal-Matrix Composites (estimated)
Major phase: Gd
Secondary phases: New unknown phase(s)
Minor phases: GdB.sub.4, GdC.sub.2, B.sub.8 C(?)
Trace phases: GdBC, Gd.sub.2 C.sub.3, Gd, B.sub.4 C, B.sub.13 C.sub.2,
GdB.sub.66, free boron(?)
TABLE II. Semi-quantitative x-ray diffraction analysis of Gd.sub.2 O.sub.3
-Gd composities.
For Cermets (measured)
Major phase: Gd.sub.2 O.sub.3 (gadolinium III oxide)
Secondary phase: Gd
Minor phases: Gd.sub.2 O.sub.3 (gadolinium oxide), GdO(?)
Trace phases: Gd.sub.2 O.sub.3 (high-temperature, metastable)
For Metal-Matrix Composites (estimated)
Major phase: Gd
Secondary/Minor phase: Gd.sub.2 O.sub.3 (gadolinium III oxide)
Trace phases: Gd.sub.2 O.sub.3 (gadolinium oxide), GdO(?)
TABLE III. Semi-quantitative x-ray diffraction analysis of B.sub.4
C-Gd.sub.2 O.sub.3 composites.
For Gd.sub.2 O.sub.3 -rich Composites (measured/estimated)
Major phase: Gd.sub.2 O.sub.3
Secondary phase: B.sub.4 C
Minor/Trace phases: GdBO.sub.3, GdB.sub.3 O.sub.6 (?), B.sub.2 O.sub.3 (?),
New unknown phase(s), other B.sub.4 C-Gd phases (?)
For B.sub.4 C-rich Composites (measured/estimated)
Major phase: B.sub.4 C
Secondary phase: Gd.sub.2 O.sub.3
Minor/Trace: New unknown phase(s), GdB.sub.4, GdC.sub.2, B.sub.8 C(?),
GdBO.sub.3, GdB.sub.3 O.sub.6 (?), B.sub.2 O.sub.3 (?), other B.sub.4 C-Gd
phases
Tables I and III indicate the confirmation of a new composition of matter.
It is noted as "New unknown phase(s)" in the tables because more than one
new phase may be present. At this time, the x-ray diffraction peaks
associated with this finding also may actually indicate (1) only a portion
of the diffraction peaks, i.e., some peaks may be masked by the presence
of larger overlapping peaks; and (2) the existence of other phases' peaks
may accidentally be included in the reported peaks below. Table IV gives
the approximate peak locations (d-spacing) and estimated peak height
(I/I.sub.o).
TABLE IV
______________________________________
Approximate d-spacing and estimated intensity of
x-ray diffraction peaks for the new phase(s)
present in B.sub.4 C--Gd and B.sub.4 C--Gd.sub.2 O.sub.3 composites.
d-spacing (Angstroms)
Intensity (I/I.sub.o)
______________________________________
8.0870 100
5.7306 .ltoreq.85
3.5660 .ltoreq.41
3.7837 .ltoreq.38
1.7299 .ltoreq.35
1.7094 .ltoreq.34
1.2942 33
2.1944 .ltoreq.33
1.9714 33
1.0262 33
1.2423 33
1.3176 32
1.7747 31
______________________________________
It is anticipated that the new phase(s) is/are ternary gadolinium
borocarbide(s).
Phases shown in Tables I-III with a (?) indicate that the existence of this
phase is likely; however, there is not enough evidence, based on the
current analysis, to say for sure if it is really present or not.
It is important to point out that subsequent heat treatment of all the
compositions of this invention will result in diminishing the major
indicated phases and promoting the growth of some of the less prevalent
phases (secondary, minor, and/or trace phases).
Examples of some refractory neutron absorbing compositions of matter, in
accordance with the invention, are given in FIGS. 7-9. FIGS. 7, 8, and 9
show the microstructures of a B.sub.4 C-Gd cermet, a B.sub.4 C-Gd.sub.2
O.sub.3 composite, and a Gd.sub.2 O.sub.3 -Gd cermet, respectively.
EXAMPLE 1
Start with -325 mesh Gd powder and -325 mesh B.sub.4 C powder. Weigh out a
60 vol. % B.sub.4 C and 40 vol. % Gd sample. Ultrasonically mix the two
powders together in a methanol slip. Filter out the methanol and dry back
the mixture. Pour the mixture into a boron nitride lined, reinforced
graphite punch and die assembly. Place assembly with mixture in a vacuum
hot press and apply 3 ksi pressure to punches. Rapidly heat to
1250.degree. C. in a vacuum of .apprxeq.10.sup.-4 torr. On passing through
1200.degree. C. increase applied pressure to 10 ksi. Hold at 1250.degree.
C. until pressing action ceases. Furnace cool part under vacuum while
maintaining 10 ksi. At 1000.degree. C. reduce applied pressure to 3 ksi
and continue cooling to room temperature. Remove part from assembly.
Result: B.sub.4 C-Gd cermet comprising
______________________________________
New unknown phase(s) 47 vol. %
GdB.sub.4 21 vol. %
GdC.sub.2 19 vol. %
other, per Table I 13 vol. %
______________________________________
EXAMPLE 2
Start with -325 mesh Gd powder and 10 .mu.m Gd.sub.2 O.sub.3 powder. Weigh
out a 75 vol. % Gd.sub.2 O.sub.3 and 25 vol. % Gd sample. Press only the
Gd.sub.2 O.sub.3 powder in a steel punch and die assembly at 10 ksi.
Remove the porous Gd.sub.2 O.sub.3 compact from the die and place it
inside the bed of the Gd powder resting in a graphite lined tungsten
crucible. Cover the crucible and place complete assembly in a vacuum
furnace. Rapidly heat to 1350.degree. C. under a vacuum of
.apprxeq.10.sup.-4 torr. Hold at 1350.degree. C. for 10 minutes to allow
molten Gd metal to infiltrate into the Gd.sub.2 O.sub.3 preform. Furnace
cool to room temperature and remove solidified melt from the crucible.
Machine away any residual Gd metal that did not infiltrate. Result:
Gd.sub.2 O.sub.3 -Gd cermet comprising
______________________________________
Gd.sub.2 O.sub.3 (III)
68 vol. %
Gd 22 vol. %
other, per Table II 10 vol. %
______________________________________
EXAMPLE 3
Start with -40 mesh Gd powder and -325 mesh B.sub.4 C powder. Weigh out a
20 vol. % B.sub.4 C and 80 vol. % Gd sample. Mechanically vibrate the
mixture for 5 minutes. Pour the mixture into a steel/graphite punch and
die assembly and compact at 20 ksi. Place the assembly with mixture into a
vacuum hot press and heat to 1200.degree. C. under a vacuum of
.apprxeq.10.sup.-3 torr. Maintain 7 ksi during heat up and hold at
1200.degree. C. until compaction ceases. Cool at 50.degree. C. per minute
under vacuum while maintaining 7 ksi applied pressure. On passing down
through 900.degree. C. reduce applied pressure to 3 ksi. Cool to room
temperature and remove billet. Extrude billet with an extrusion ratio of
1:10 at a temperature of 980.degree. C. Finally, heat treat extruded part
for 3 hours at 1100.degree. C. Result: B.sub.4 C-Gd metal-matrix composite
comprising
______________________________________
Gd 55 vol. %
New unknown phase(s) 28 vol. %
GdB.sub.4 8 vol. %
GdC.sub.2 6 vol. %
other, per Table I 3 vol. %
______________________________________
EXAMPLE 4
Start with -40 mesh Gd powder and 10 .mu.m Gd.sub.2 O.sub.3 powder. Weigh
out a 15 vol. % Gd.sub.2 O.sub.3 and 85 vol. % Gd sample. V-blend the
powders together for 30 minutes. Place the mixture in a susceptor crucible
and place in an induction melting furnace. Inductively heat the assembly
at 1400.degree. C. in a vacuum of 10.sup.-2 torr for 1 hour. Furnace cool
to room temperature and remove billet from crucible. Extrude billet at an
extrusion ratio of 1:8 at a temperature of 900.degree. C. Finally, heat
treat the part at 1000.degree. C. for 2 hours. Result: Gd.sub.2 O.sub.3
-Gd metal-matrix composite comprising
______________________________________
Gd 64 vol. %
Gd.sub.2 O.sub.3 (III)
25 vol. %
other, per Table II 11 vol. %
______________________________________
EXAMPLE 5
Start with -20 .mu.m B.sub.4 C powder and -10 .mu.m Gd.sub.2 O.sub.3
powder. Weigh out a 65 vol. % B.sub.4 C and 35 vol. % Gd.sub.2 O.sub.3
sample. Prepare a sterically stabilized colloidal suspension and
ultrasonicate the slip for 3 minutes. Pressure cast the slip at 80 psi for
24 hours. Remove green part from the pressure caster and place in a
graphite punch and die assembly. Place the assembly in a hot press. Hot
press the green body at 1950.degree. C. with an applied pressure of 5 ksi
in flowing argon. Cool at 25.degree. C. per minute down through
500.degree. C. while maintaining 5 ksi applied pressure. On passing
500.degree. C. cool at 100.degree. C. per minute while maintaining only 3
ksi applied pressure. Remove part from die at room temperature. Result:
B.sub.4 C-rich B.sub.4 C-Gd.sub.2 O.sub.3 composite comprising
______________________________________
B.sub.4 C 62 vol. %
Gd.sub.2 O.sub.3 33 vol. %
other, per Table III 5 vol. %
______________________________________
EXAMPLE 6
Start with 1-3 .mu.m B.sub.4 C powder and -10 .mu.m Gd.sub.2 O.sub.3
powder. Weigh out a 30 vol. % B.sub.4 C and 70 vol. % Gd.sub.2 O.sub.3
sample. Apply electrostatic dispersion techniques to obtain a slip.
Ultrasonicate slip for 5 minutes in an ice bath. Slip cast slurry into
plaster mold. Remove green part and vacuum dry at 100.degree. C. and 50
millitorr for 12 hours. Place green part in inert gas furnace and fire at
1850.degree. C. for 1 hour. Cool at 50.degree. C. per minute to room
temperature. Result: Gd.sub.2 O.sub.3 -rich B.sub.4 C-Gd.sub.2 O.sub.3
composite comprising
______________________________________
Gd.sub.2 O.sub.3 66 vol. %
B.sub.4 C 27 vol. %
other, per Table III 7 vol. %
______________________________________
OTHER APPLICABLE COMPOSITES
Although the best compositions for neutron absorption are those using Gd
metal (or alloy), upon which this invention is based, there are several
other metals that may be combined with B.sub.4 C or Gd.sub.2 O.sub.3 to
form other cermet and metal-matrix compositions, which may be useful for
the same or other applications. The metals that can be used in accordance
with the methods of this invention are Cd, In, Te, Pb, Ce, Pr, Nd, Sm, Eu,
Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu, and alloys thereof.
Changes and modifications in the specifically described embodiments can be
carried out without departing from the scope of the invention which is
intended to be limited only by the scope of the appended claims.
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